Method for producing ultra fine-grained microstructure in ferrous alloys

ABSTRACT

Ferrous alloys containing up to about 25 percent Ni, 15 percent Mn, 1.2 percent C, 0.5 percent N are cooled to form a martensitic or bainitic structure. The material is then worked to an extent sufficient to remove the principal nucleating effect of the prior austenite boundaries and other microstructural interfaces, so that when subsequently heated into the multiphase region, recrystallization occurs by random nucleation of extremely fine austenite crystals thoughout the material. Exceptionally fine equiaxed grains in the micron and submicron range are achieved, thereby providing unique combinations of both increased strength along with increased ductility and increased notch toughness.

United States Patent [1 1 Miller METHOD FOR PRODUCING ULTRA FINE-GRAINED MICROSTRUCTURE 1N FERROUS ALLOYS 2] App1.No.: 182,331

[52] 11.5. C1. 148/12, 148/124 [51] Int. Cl. C21d 7/14, C21d 9/46 [58] Field of Search 148/12, 12.4, 143

[56] References Cited UNITED STATES PATENTS 2/1966 Grange et a1. 148/12.4

3/1970 Grange 148/124 1451 Aug. 28, 1973 3,196,053 7/1965 Hodge 148/124 Primary ExaminerW. W. Stallard Attorney-Arthur J. Greif [57] ABSTRACT Ferrous alloys containing up to about 25 percent Ni, 15 percent Mn, 1.2 percent C, 0.5 percent N are cooled to form a martensitic or bainitic structure. The material is then worked to an extent sufficient to remove the principal nucleating effect of the prior austenite boundaries and other microstructural interfaces, so that when subsequently heated into the multiphase region, recrystallization occurs by random nucleation of extremely fine austenite crystals thoughout the material. Excep tionally fine equiaxed grains in the micron and submicron range are achieved, thereby providing unique combinations of both increased strength along with increased ductility and increased notch toughness.

14 Claims, 9 Drawing Figures TEMPERATURE, F.

l I l l 1 \O g 40 400 3 Q Q r 30- 360 3 "U k E- i 1% Lu 20 320 q k I b q O Q HARD/V555 TEMPERATURE, 0.

Patented Aug. 28, 1973 AUSTE/V/TE', VOLUME 8 Sheets-Sheet 1 TEMPE/PA 70/95, "F. I200 HARD/V585 l 1 l 1 I 1 l a D HARD/V588, DPH

TEMPE/W1 TURE, C.

Patented Aug. 28, 1973 8 Sheets-Sheet 2 FIG. 2.

(b) /0,000 X ANNEALED /6 Hrs. 6000. (25 Tl It} I900 X M) /5,300X

COLD WORKED AND ANNEALED /6 Hrs. 60096. (29% T) (E/MZLUUHO x WARM WORKED AT.6000. (9-2196 T) AND 6400. (f-30 967) Patented Aug. 28, 1973 8 Sheets-Sheet 5 r H l m I l 1 1 H n 1 I I p 0 d, a. M A E 5 M r. G .14 W m ID H I 1 m 0 WM. N mu CA I .m oa m M 6 p v 0 0 w 4 2 o TIME A T 600C.

0 COLD WORKED AND ANNEALED U ANNEALED TEST TEMPERATURE, F.

Patented Aug. 28, 1973 3,755,004

8 Sheets-Sheet 6 0 com won/(0 AND ANNEALED ,D n90 0 WA RM WORKED mo TENS/LE STRENGTH ksi Y/ELD STRENGTH IOO- ksi

0 REDUCTION OF AREA D U ELO/VGAT/OIV 20- AUSTE/V/ TE CONTENT l l I l l l l l TEMPERATURE, C.

Patented Aug. 28, 1973 3,755,004

8 Sheets-Shoot 7 (b) /6 Hrs.

(0) I00 Hrs. (d) 300 Hrs.

MICROSTRUCTURE 0F F's-20% Ni ALLOY COLD WORKED AND HEATED 7'0 500C F0 VARIOUS TIMES Patented Aug. 28, 1973 8 Sheets-Shoot 8 PRODUCT THICKNESS, INCHES 0.4 0.3

l l l l 30 REDUCTION //V THICKNESS B) COLQROLL/NG,

FIG. 8.

METHOD FOR PRODUCING ULTRA FINE-GRAINEI) MICROSTRUCTURE IN FERROUS ALLOYS This invention pertains to a method of imparting an ultrafine-grained microstructure to steel alloys containing martensitic or bainitic constituents and to those which are capable of conversion to martensite or bainite. The alloys processed by the method of this invention manifest a unique combination of high ductility and notch toughness together with high strength.

The desirable effect of extremely fine grain size on mechanical properties has long been recognized. However, in commercial practice, desired mechanical properties have generally been achieved through the utilization of alloying additions. Due to ever-increasing cost of alloying elements, however, the use of heat treatments to effect significant decreases in grain size so as to achieve such desirable mechanical properties has only come to the forefront in recent years. Thus, U.S. Pat. No. 3,178,324, by employing a rapid, cyclic heat treatment, describes the production of grains as small as 2.5 m (ASTM No. 14); thus, eliminating the need for expensive vanadium and columbium. In US. Pat. No. 3,444,01 l,- the need for costly nickel to achieve low temperature toughness is lessened, in part, by the utilization of a tempering heat treatment to achieve a dispersion of fine austenite grains (part or all of which is converted to martensite on cooling) in a ferrite matrix.

It has now been discovered that if ferrous compositions capable of conversion to structures such as martensitic platelets, or bainitic, ferrite-carbide aggregates are (l) austenitized and (2) cooled to produce such martensitic or bainitic structures, and (3) then worked to a sufficient degree (as hereinafter described) followed by (4) heating in the multiphase region, wherein one of the phases is austenite; microstructures are achievable with grain sizes in the micron and submicron range, i.e., smaller than ASTM No. 16.

It is therefore a principal object of this invention to provide a thermomechanical method for achieving an ultra fine-grained microstructure in a wide variety of ferrous alloys.

The foregoing and other objects will become more apparent from the following detailed description, when read in conjunction with the appended claims and Figures in which:

FIG. 1 is a plot of the room temperature hardness and the austenite content of an Fe5.7 percent Mn-0.l percent C alloy after 1 hour at temperature;

FIG. 2, a through f, are electron transmission micrographs showing the effect of severe working prior to or during annealing in the critical multiphase range;

FIG. 3 depicts the effects of severe cold working on mechanical properties after annealing at 600 C;

FIG. 4 depicts the effects of severe cold working on mechanical properties after annealing at 640 C;

FIGS. 5a and 5b show the effects of severe cold work on toughness (one-half width CVN specimens);

FIG. 6 depicts the effects of severe working at temperatures within the critical multiphase region, on mechanical properties;

FIGS. 7a, 7b, 7c and 7d are micrographs showing the long time stability of Fepercent Ni alloys, cold worked and tempered at 500 C; and

FIG. 8 is a plot of the proportion of equiaxed, finegrained structure in the cross section of initially 1% inch thick plate, cold rolled to varying degrees and annealed. Micrographs at 21 ,OOOX show the structure at the center of the cross section.

In its broadest aspects, the method of this invention is applicable to any steel capable of conversion to martensite or bainite, when cooled from an austenitizing temperature. If such an alloy were then tempered in the multiphase region (i.e., at a temperature above that at which austenite formation begins), austenite would nucleate principally in boundary regions, such as prior austenite grain boundaries, martensitic plate interfaces (or bainitic, aggregate interfaces) and perhaps low angle subgrain boundaries. Thus, even after long periods of annealing within this region, the microstructure would retain an appearance similar to that of the original structure. However, if these materials are deformed, to a critical extent, the relative nucleating effect of such boundary regions may be greatly diminished and nucleation would occur randomly. This random nucleation provides an entirely different morphology, that of an equiaxed, ultra fine-grained microstructure. In a more specific embodiment, the inventive method is applicable to those compositions which contain a sufficient amount of the austenitizing elements, e.g. Ni, Mn, N and C, so that the austenite which forms during the aforesaid tempering treatment is retained on cooling to room temperature. In another embodiment, these austenitizing elements are insufficient to impart such stability, and part or all of this austenite isconverted, on cooling, to well-known transformation products such as pearlite, bainite or martensite. In a further embodiment, in which martensite is produced during such cooling, the material is further tempered at a lower temperature to provide increased ductility.

Since practical, commercial heat treatments often differ markedly from equilibrium conditions and since the extent of the multiphase phase regions (ferrite austenite carbides) for alloys such as Fe-Mn-C, Fe- Ni-C and Fe-Ni-Mn-C is not precisely known, the tempering range must be defined in terms of nonequilibrium designations. For purposes of this invention, A," is that temperature at which the start of austenite nucleation may be detected by holding at temperature for one hour. A," is that temperature at which the finish of the transformation to austenite is detected, for a one-hour hold at temperature. Analogous, A -A," regions have been previously determined by various investigators employing metallographic and hardness measurements. In dealing with many of the compositions to which the instant method is applicable, reliance on such measurements only, may provide erroneous results. An illustration of this problem may be found in the following example. Specimens of a 5.7 percent Mn, 0.1 percent C steel were heated for 1 hour at temperatures in the range of 520 to 720 C. Since the presence of austenite was not detected by light microscopy, X-ray diffraction was employed to determine quantitatively the volume fraction of austenite after cooling to room temperature. These data and the corresponding hardness data are shown in FIG. I. Austenite formed at all temperatures investigated. If hardness measurements alone were performed, it might have been concluded (erroneously) that the A,-A, range extended from about 650 C to 740 C, since the austenite which forms below 650 C does not transform and hence produces a slight softening rather than hardening. This austenite stability (below 650 C) is believed to be due to the combined effect of smaller crystal size (in comparison to the austenite grain size which results at higher temperatures) and enrichment of the austenite in Mn and C due to partitioning (with the concurrent depletion of Mn and C in the ferrite). It may be seen, therefore, that it is desirable to know the A, value of a particular composition with some certainty, so that the tempering treatment of this invention may preferably be performed in the lower portion of A,-A, range, e.g. within the range A, 150 C. Such lower temperature tempering, will provide the advantage of even smaller grains, (and therefore even greater improvement in mechanical properties) when compared with that of tempering near the upper portion of A,,-A, range.

The enhanced effect of severe working (well above the critical minimum) prior to or during tempering in the critical range may be seen by reference to FIGS. 2 through 6. The treatments represented therein were performed on a Fe-5.7 percent Mn-0.l percent C steel.

FIGS. 2a and 2b show typical micrographs of the annealed martensitic structure, without any prior working. FIG. 2a shows relatively large austenite crystals in a prior austenite grain boundary, and thin, elongated austenite crystals in the boundaries between martensitic plates (the dark streaks in 2b). Even after relatively long annealing, the structure retains an appearance resembling the original martensite.

Cold working (by swaging to 85 percent reduction in cross-section) prior to annealing, removes the effect of these boundaries as the principle nucleation source, and recrystallization occurs by random nucleation of ultrafine austenite crystals as shown in FIGS. 2c and 2d. The grain size depicted herein is 0.5 pm (ASTM No. 19

'l he effect on microstructure of working within the critical range may be seen in FIGS. 2e and 2f. Specimens were warm worked by rolling out ofa molten lead bath to 75 percent reduction in thickness. Multiple passes were employed in which the specimens were returned to the lead bath between each pass. It may be seen that the microstructure resembles that of FIGS. 2c and 2d, except that it is not fully annealed. The grain size is even smaller than that of the cold worked and annealed structure (2d).

The influence of the foregoing microstructures on mechanical properties is shown in FIGS. 3 through 6. FIGS. 3 and 4 show the amount of austenite formed and the effect of percent prior cold reduction on 1 tensile properties after tempering at 600 and 640 C, respectively. It may be noted that although the austenite contents are quite similar, the material which has been severely worked prior to annealing shows an improvement in nearly all properties. Note especially, that yield strength increases of about 30 percent were achieved, with no sacrifice (improvement in most cases) in ductility, as measured by reduction in area and elongation. The ultra fine-grained microstructure is particularly noteworthy in its effect on toughness, as shown in FIGS. 5a and 5b. A significant improvement in notch toughness (as measured by ft.-lbs. of energy absorbed) is achieved, particularly at low test temperatures.

In FIG. 6, results from specimens cold worked 60 percent and tempered at four temperatures are compared with results from specimens warm worked percent in the critical tempering range. Since, in both cases, the extent of deformation was above the critical amount (i.e., that amount which substantially negates the nucleating effect of boundaries) it may be seen that where the tempering temperatures are the same, quite similar properties are obtained. This similarity results from the fact that, as previously noted in FIG. 2, the microstructures are also quite similar.

Corresponding results were also obtained for various Fe-Ni alloys. To evidence the long time, high temperature stability achieved by the instant method, Fe-20 percent Ni specimens were cold worked by rolling to an percent reduction in thickness, prior to tempering in the two-phase region. As shown in the electron trans mission micrographs of FIG. 711, after 30 minutes at 500 C, the crystal size is considerably less than 0.25 am. After 16 hours at temperature, the alloy has re crystallized to uniform, equiaxed grains of about 0.3 nm (FIG. 7b). After 100 hours (70) the grain size is about 0.4 m and even after 300 hours it is only about 0.6 nm (7d). In specimens held as long as 2000 hours at 500 C, the grain size was still less than 1 um. Mechanical properties of five different Fe-Ni-C compositions are listed in Table I, along with the corresponding heat treatment and resulting grain size. In addition to the substantial improvement in mechanical properties evidenced, these latter results also depict the exceptionally long time thermal stability of alloys produced by the method of this invention.

PROPERTIES OF FE-NI ALLOYS Austen- Grain Upper Lower ltv Ferrite size yield yield Treatment, volume, volume, (1, stress, stress,

Alloy hours] 0. percent percent nn p.s.i. p.s.i.

, 300 500 40 .40 122, 000 10.1, 500 1, 000/500 54 40 07 101,000 0x, 500 1, 000/450 32 (i8 37 137, 500 I25, 000

Q 3 1500 40 51 .55 1015,1100 0x, 100 I 1, 000/500 51 411 .73 00,000 00, 000 1, 000/450 37 (i3 42 128, 000 111i, 500

v- 3 .215 74 50 112, 700 101, .200 Ff 0053C 1, 000/500 31 00 .00 105, 000 00, 000 1, 000/450 15 51 125, 000 111, 300

The critical minimum amount of working required to produce the desired random nucleation will depend on a variety of factors. Thus, the degree of deformation required, will depend upon the initial structure, eg martensite or bainite. Thicker specimens will, in general require a greater amount of deformation to provide an equiaxed microstructure, which is essentially uniform across the entire cross-section (as compared with a structure in which only the surface region has received the required amount of deformation). The type of deformation employed (rolling, swaging, drawing, etc.) will itself play a role, due to the differing effects of these modes of deformation on the internal energy of the metal. Even for a specific type of deformation, e.g. by rolling, the amount of deformation will depend, to some extent, on the roll diameter, the surface conditions of the rolls, the type of lubrication and the amount of reduction per pass.

To achieve the desired grain refinement, the extent of working (whether at room temperature or above) must be such that the multi-phase region tempering will result in a substantially equiaxed structure. This critical amount may be determined, for a particular set of conditions, by a rather simple test, such as that outlined below.

Starting with A inch-thick plate of the Fe-5.7 percent Mn-0.l percent C steel, specimens were cold rolled to a reduction in thickness varying from percent to 70 percent; and then annealed for 16 hours at 600 C. When the amount of deformation is small, only the region just below the surface is cold worked sufficiently to produce the desired equiaxed, fine-grained microstructure. As the amount of cold reduction is increased, this region extends further into the center of the piece, until at a critical value, the microstructure is essentially uniform across the entire thickness of the piece. While the ultra fine equiaxed grains are too small to be delineated by an ordinary light microscope, this instrument is nevertheless capable of differentiating between the equiaxed surface region and the center region which has not been deformed sufficiently. Thus, the above annealed specimens were observed and measured under ordinary light at lOOOX. The ratio in the thickness of the equiaxed recrystallized region to the total thickness was plotted (FIG. 8) against the amount of deformation each specimen had received. For the particular conditions of deformation in this experiment, it may be seen that at about 40 percent reduction in thickness, substantially the whole plate, has achieved the ultrafinegrained structure. Electron transmission foils were also prepared from the center of the plate at these various deformations. The abrupt change in nucleation morphology at about 35 percent reduction is clearly evident. The effect of deformation well above the critical value, in effecting even further grain refinement, may be seen in the micrograph at 70 percent reduction.

As stated previously, the instant method may be effectually applied to all compositions, which on cooling can be transformed to martensite or bainite. Thus, for example, in a composition containing a level of nickel far in excess of about 25 percent the austenite will be too stable for such a necessary conversion. Similarly, a level of manganese far in excess of about percent (or a lesser combination of manganese, nickel and/or nitrogen) will provide a composition which is too stable to provide the desired starting structure of martensite or bainite. In addition to austenite stability, other compositional constraints may be noted. Generally, an alloy containing greater than about 1.2 percent carbon or about 0.5 percent nitrogen will be too hard to be worked sufficiently, such that the recrystallization during tempering in the multiphase region will result in the desired equiaxed morphology.

In the specific examples provided for illustration of the invention, the quantity of austenitizing elements (Mn, Ni and C) was sufficient, so that the ultrafine austenite which formed on tempering in the critical range, was stable even after cooling to well below room temperatures. lf leaner compositions are employed (e.g., Fe-2 percent Mn-l.0 percent Ni-O.2 percent C) in which the austenite will be comparatively less stable, the austenite can then transform (depending on cooling rates employed) to the well-known transformation products. Thus, by utilization of continuous slow cooling or by isothermal holding at relatively high temperatures, pearlite and feathery bainite will be produced, whereas faster cooling rates would provide acicular bainite or martensite. In the latter instance, undesirable embrittlement due to the martensite could be obviated by well-known treatments such as interrupted quenching (martempering) or by tempering the martensite.

I claim:

1. A method for treating ferrous structures selected from the group consisting of martensite or bainite so as to achieve an essentially uniform, ultrafine grain size, which comprises;

a. deforming the structure to an extent sufficient to substantially diminish the nucleating effect of the prior boundary regions;

b. tempering the deformed ferrous structure to a temperature above the A," and below the A, for a time sufficient to effect the growth of substantially equiaxed, ultrafine grained austenite crystals; and

c. cooling the so-worked alloy to below the A," temperature.

2. The method of claim 1 wherein said structure is martensite.

3. A method for producing an essentially uniform, ultrafine grain size in ferrous alloys capable of conversion to martensite or bainite, which comprises;

a. heating said alloy to a temperature within the austenitizing range for a time sufficient to effect a sub stantially complete transformation to austenite;

b. cooling the so-formed austenite to form a microstructure consisting substantially of martensite or bainite;

c. deforming the structure of (b) to an extent sufficient to substantially diminish the nucleating effect of the prior boundary regions;

d. tempering the so-worked alloy to a temperature above the A," and below the A," for a time sufficient to effect the growth of equiaxed ultrafine grained austenite crystals; and

e. cooling the so-worked alloy to below the A," temperatures.

4. The method of claim 3 in which the extent of said deformation is such that the recrystallization effected during said tempering results in a substantially uniform equiaxed ultrafine grain microstructure.

5. The method of claim 4 wherein said alloy is selected from the range, consisting essentially of the austenitizing elements of up to about 25 percent nickel, up to about 15 percent manganese, up to about 1.2 percent carbon, up to about 0.5 percent nitrogen and mixtures thereof with the balance iron and incidental steelmaking impurities.

6. The method of claim in which said alloy contains a sufficient amount of said austenitizing elements to provide sufficient austenite stability so that the austenite so formed during said tempering is retained during said cooling below the A," temperature.

7. The method of claim 5 wherein the amount of said austenitizing elements is insufficient to provide sufficient stability to the austenite formed during said tempering, whereby the austenite transforms to decomposition products selected from the group consisting of pearlite, martensite and bainite.

8. The method of claim 7 wherein said cooling is at a rate sufficiently fast to produce martensite.

9. The method of claim 8 wherein said martensite is retempered at a temperature below the A,

10. The method of claim 7 wherein said cooling is interrupted to convert said equiaxed, ultrafine grained austenite crystals to a martempered structure.

11. The method of claim 4 wherein said deformation is performed at a temperature above the A," and below the A,.

12. The method of claim 4 wherein said deformation is performed at a temperature below the A,."

13. The method of claim 3 wherein said tempering is conducted at a temperature within the range A, to A, plus C.

14. The method of claim 4 wherein said deformation is at least that required to increase the internal energy of the alloy to that represented by the increase achieved by a reduction in thickness of 35 percent, for a r& inch thick plate. 

2. The method of claim 1 wherein said structure is martensite.
 3. A method for producing an essentially uniform, ultrafine grain size in ferrous alloys capable of conversion to martensite or bainite, which comprises; a. heating said alloy to a temperature within the austenitizing range for a time sufficient to effect a substantially complete transformation to austenite; b. cooling the so-formed austenite to form a microstructure consisting substantially of martensite or bainiTe; c. deforming the structure of (b) to an extent sufficient to substantially diminish the nucleating effect of the prior boundary regions; d. tempering the so-worked alloy to a temperature above the ''''As'''' and below the ''''Af'''' for a time sufficient to effect the growth of equiaxed ultrafine grained austenite crystals; and e. cooling the so-worked alloy to below the ''''As'''' temperatures.
 4. The method of claim 3 in which the extent of said deformation is such that the recrystallization effected during said tempering results in a substantially uniform equiaxed ultrafine grain microstructure.
 5. The method of claim 4 wherein said alloy is selected from the range, consisting essentially of the austenitizing elements of up to about 25 percent nickel, up to about 15 percent manganese, up to about 1.2 percent carbon, up to about 0.5 percent nitrogen and mixtures thereof with the balance iron and incidental steelmaking impurities.
 6. The method of claim 5 in which said alloy contains a sufficient amount of said austenitizing elements to provide sufficient austenite stability so that the austenite so formed during said tempering is retained during said cooling below the ''''As'''' temperature.
 7. The method of claim 5 wherein the amount of said austenitizing elements is insufficient to provide sufficient stability to the austenite formed during said tempering, whereby the austenite transforms to decomposition products selected from the group consisting of pearlite, martensite and bainite.
 8. The method of claim 7 wherein said cooling is at a rate sufficiently fast to produce martensite.
 9. The method of claim 8 wherein said martensite is retempered at a temperature below the ''''As.''''
 10. The method of claim 7 wherein said cooling is interrupted to convert said equiaxed, ultrafine grained austenite crystals to a martempered structure.
 11. The method of claim 4 wherein said deformation is performed at a temperature above the ''''As'''' and below the ''''Af.''''
 12. The method of claim 4 wherein said deformation is performed at a temperature below the ''''As.''''
 13. The method of claim 3 wherein said tempering is conducted at a temperature within the range ''''As'''' to ''''As plus 150* C''''.
 14. The method of claim 4 wherein said deformation is at least that required to increase the internal energy of the alloy to that represented by the increase achieved by a reduction in thickness of 35 percent, for a 1/2 inch thick plate. 